Castable heat resisting iron alloy



Jan. 12, 1965 A. ROY ETAL 3,165,400

CASTABLE HEAT RESISTING IRON ALLOY Filed June 2'7, 1961 8 Sheets-Sheet l INVENTORS. 17716166 74% )n d/ier Z. J'omzrrg.

A TTORNE Y6 ,Jan. 12, 1965 A. ROY ET-AL 3,

'CASTABLE HEAT RESISTING IRON ALLOY Filed June 27, 1961 8 Sheets-Sheet 2 IN VENT OBS. fimeie e 770 M d/P'ef Z, J'amrrzy ATTORNEYS.

Jan. 12, 1965 A. ROY ETAL CASTABLE HEAT RESISTING IRON ALLOY Filed Jun 27, 1961 8 SheetsSheet 3 INVENTORS, flmeine 70y BY 14/4 /767 J, 1.707772?? Wit/441M ATTORNEYS" Jan. 12, 1965 A. ROY ETAL 3,165,400

CASTABLE HEAT RESISTING IRON ALLOY Filed June 27, 1961 8 Sheets-Sheet 4 IN VEN TORS A TTORNEYS.

Jan. 12, 1965 A. ROY ETAL 3,165,400

CASTABLE HEAT RESISTING IRON ALLOY Filed June 27, 1961 8 Sheets-Sheet 5 .BY I

A TTORNEYS. I

Jan. 12, 1965 A. ROY ETAL 3,165,400

CASTABLE HEAT RESISTING IRON ALLOY Filed June 27, 1961 8 Sheets-Sheet 6 INVENTOHS. flmeaiie 0y I M/d/Zer I 3277-1272? M a: A a/w A TTORNEYS.

Jan. 12, 1965 A. ROY ETAL 3,155,400

CAS'I'ABLE HEAT RESISTING IRON ALLOY Filed June 27, 1961 8 Sheets-Sheet 7 AT T ORNE Y-S A. ROY ETAL CASTABLE HEAT RESISTING IRON ALLOY Jan. 12, 1965 8 Sheets-Sheet 8 Filed June 27, 1961 United States Patent 3,165,400 CASTABLE HEAT RESISTING HRON ALLOY Amedee Roy, Ferndale, and Walter E. Iiominy, Detroit, Mich, assignors to Chrysler Corporation, Highland Park, Mich, a corporation of Delaware Filed June 27, 1961, Ser. No. 119,902

21 Claims. 1 (Cl. 75-426) This invention relates to iron base alloys austenitic at room and operating temperatures and possessing high stress rupture properties at temperatures above 1200 F. and up to 1500. F. and in'excess thereof even up to 1700 F. without the use of substantial amounts of the strategic elements nickel (Ni) and cobalt (Co) and even in the substantial absence of one or both. It especially concerns alloys of this character possessing in high temperature service stainless properties for resisting oxidation and providing substantial strength and load carrying ability and great resistance to creep and deterioration by heat. It also particularly refers to iron base alloys of the aforesaid character useful for producing cast articles of complicated shapes for instance, by precision casting processes. Moreover, it has specific reference to alloys of the foregoing character and properties which are relatively high in carbon content and which are weldable by known procedures.

The mass production of many mechanical parts depends upon the development of low cost, non-strategic alloys which provide adequate strength and oxidation resistance at the high temperature service conditions at which such parts are to operate. turbine parts, internal combustion engine valve parts and heat treat fixtures. Iron base alloys high in carbon and ferritic at room temperatures are not suitable for applications such as these where temperatures in the range of 1200 F. to 1'700 F. and relatively high stress conditions are encountered. Moreover, they are either not suitable for welding or may only be welded with substantial difiiculty.

Cobalt base alloys'rich in nickel, for example Stellite alloy H831 and nickel base alloys usually referred to as highly alloyed super alloys, have been proposed for meeting some of the properties required for such rigorous service but these high strategic element alloys cannot at present be considered for most mass produced structures because of their high cost and the high content of strategic elements such as cobalt and/or nickel they necessarily contain. Moreover, attempts made to produce substitute castable alloys low in the strategically scarce elements cobalt and/or nickel, have not insofar as we are aware, produced any having the foregoing physical properties required at temperature between 1200" to1700 F. nor have such attempts produced any alloy approaching the service life possible with available super alloys predominant in cobalt and nickel. Even heat treatment of Typical examples are furnace and such alloys has not developed satisfactory high temperscarce elements is the necessity'in many cases of employ-. ing special high temperature heat treatments in order to develop their high temperature properties. This is very costly. Moreover, even with these heat treatments We are not aware that any ferrous base alloy for the above 3,165,46 Patented Jan. 12, 1965 ice application has been commercially available providing a hour rupture strength of at least 20,000 p.s.i.

The present invention aims to solve the aforesaid problems and provide suitable iron base alloys not requiring large amounts of any strategic elements especially cobalt and/ or nickel and capable of providing the required high physical and service properties. It is based upon the discovery that iron base alloys are structure sensitive at high temperatures and that obtaining desired room temperature and high temperature properties aforesaid is dependent upon the proper proportioning in low critical amount of certain austenite forming elements especially carbon and nickel and NaCl-type carbide formers in association with other essential elements of the composition.

More particularly it is based upon the discovery that improved room temperature and high temperature properties are imparted to iron base alloys for casting into complex shapes, by the presence of substantial amounts of carbon and by the presence of multiple and controlled additions of the NaCl-type of cubic carbide-forming elements in a stable and balanced austenitic composition containing critical amounts of carbon and nickel and that the NaCl-type of cubic carbides of a selected group of elements are found to possess a very high degree of stability at elevated temperatures conducive to the making of low nickel, cast austenitic stainless steels which may be co-- balt free, with outstanding stress rupture properties at high temperatures of 1500 F. and even higher. In particular cases by a proper combination of critical amounts of certain carbide forming elements with nitrogen, the alloy may even be made nickel free.

A fundamental object is therefore to provide new and improved heat resisting castable iron base alloys austenitic at room temperatures, relatively high in carbon, low in A further object is to provide new and improved high temperature iron base alloys as in the preceding objects which do not require special heat treatment to develop vastly improved high temperature properties but which may in certain cases by still further improved thereby.

A particular object is to provide high carbon stainless iron base alloys austenitic at room temperature and low in strategic elements which possess at high temperatures of 1500 F. properties comparable or superior to those attained by so-called super alloys high in one or more of the strategic elements nickel and/ or cobalt.

A specific object is to provide a heat resisting stainless iron base alloy containing in the order of 60% and more iron by Weight and which may be low in nickel and 'free of cobalt, which alloy is austenitic at room temperature, suitable for producing cast turbine blades and the like, and will provide articles of this character with adequate ductility and strength at room temperature and the ability to resist creep, oxidation and scaling at temperatures of 1500 F. and greater encountered in the type of service to which such articles are subject.

Another specific object is to provide an alloy composition comprising a wholly austenitic iron, solid solution matrix relatively high in carbon content in which insoluble primary and lamellar-type carbides are embedded.

and which contains dot-type carbides and angular intermetallic'compounds randomly distributed in the matrix and semi-continuously in the grain boundaries.

A further object is to produce an iron base alloyfcomposition relatively high in carbon content providing a V izedby the absence of: secondary precipitates except for a few localized secondary carbides at precipitated from theaustenitic matrix. v p FIGURE .2 represents a photomicrograph' (1000 of a cast alloy No. .123 containing 0.54% carbon, which alloy is also not within the scope of our invention and matrix having a stable austenitic face-centered cubic structure (FCC) as lean as possible in the requisite nonstrengthening substituent type austenitic forming elements 'in order to facilitate the addition of optimum amounts of matrix strengtheners Without overstraining the iron lattice of the structure at both room and elevated temperatures. v t

A particular object is to produce an iron base alloy composition low in strategic elements and relatively high in carbon that has good physical properties at high temperaturesand that is readily weldable.

"It is a further object toprovide an alloy matrix Whose 1 composition comprises an iron base alloy relatively high.

in carbon content containing dissolved and combined nitrogen, in controlled concentration and balanced to the total amount of nitride formers such as chromium, manganese and columbium present therein.

A further specific object is to provide .a low cost alloy metal having optimum load-carrying ability and creep which was taken from a /4,'. diameter test bar heated for about 100 hours at 1500F. The alloy' is austenitic. The light colored area indicated by the letterzA signifies the austenitic phase, the letters B represent the eutectic (ledeburite) phases .randomlydispersed throughout the structure, and .the letter C refers to dot-like'secondary carbides randomly dispersed within the matrix and locally concentrated toward the eutectic phases.

FIGURE 3 represents a photomicrograph(l000 of a cast alloy similar to No. 123 but containing 0.7% carbon taken from 2; Mi diameter test bar heated for about 100 hours at 150i) F. The alloy is austenitic.

resistance at high temperatures between 1200 F. to

1700 F. which comprises a stable iron base castable alloy austenitic at room and operating temperatures including as essential ingredients, predetermined relatively small amounts of austenite forming elements from the group consisting of carbon, nickel and manganese, such consisting at least of carbon in substantial amount and of nickel and predetermined relatively small amounts of a pluralityof carbide-and/ or nitride forming elements including molybdenum, tungsten, and at least one selected from the group consisting of columbium, tantalum,

. the preceding objects that includes, free or uncombined chromium in amount sur'ficient to render the metal stainless in character and oxidation resistant at hightemperatures between 1200? F. to 1800 F An additional object is to provide an ironibasealloy having high stress .resistance properties and a face-centered cubicsolid solution 'structure'stable .at room temperature that includes in addition to chromium and criticalamounts of carbon, small predetermined amounts of structures.

It is similar in general structural characteristics to the alloy=of FIGURE 2 but differs in that there is. a greater amount of randomly dispersed eutectic phases'B'in the austenitic phase A. Also, there is a greater'concentra tion of dot-like secondary carbides C within the matrix.

FIGURE .4 represents aphotomicrograph (lOOOX) of a cast alloy No. 120 of our invention containing 0.83%

carbontaken from a' diameter test bar heated for about 100 hours at 150031. The alloy is austenitic. The structure is significantly different from those of FIGURES l to 3, inclusive, in that the eutectic (ledeburite) phases B form a distinctive well defined network surrounding the austenitic phases A imbedding dot-like secondary can bides C which are heavily concentrated toWard'the-net-" work and more uniformly distributed throughout the austenitic phase A than in the previously described micro- FIGUREo' represents a photomicrograph (IOOOX) of a cast alloy No. 6 of our invention containing 1;05% carbon taken from-a Mi" diameter test'bar heated for about 100 hours at 1500 F. The alloy is austenitic andis of Y the general character of that of FIGURE 4, butis notable by its heavier well, defined. network of the eutectic (ledeburite) phases. B outlining the austenitic matrix areas A and heavier concentrates'of fine secondary carbide precipitates C within'the austenitic phase.

FIGURES 6 and 7' represent photomicrographs' to that of FIGURES 4 and 5, but are characterized by the presence of increasing density in the eutectic (ledeburite) network. a

FIGURE 8 is a graphshowing the efiect of varying the carbon contentof-a basic alloy composition such as alloy nickel and in particular instances, nickel and/or man- 1 ganese .for stabilizing the austenitic structure by substitution for iron, and other strengthening elements such as nitrogen, molybdenum, tungsten, columbium, tantalum,

titanium, and vanadium in predetermined combinations in solid solution or asa precipitate for the purpose of matrix strengthening-without overstraining the iron lattice.

Other objects and advantages of our invention will appear from the following additional description and from the drawings wherein:

FIGURE 1 represents a photomicrograph (1000 of a cast alloy No. 124 containing 0.38% carbon, which I alloyis not within the scope of our invention and which Was takenfrom a A" diameter test bar heated for about hoursat 1500 F. .The alloy is austenitic. The

light colored area indicatedby letter A represents the austenitic phase, the black areas indicated by theletters B represent eutectic (ledeburite) phases randomly dispersed throughout the structure. The alloy is characterslowly to No. 6 on the 100 hour stress rupture strengthat 1500 F. and room temperature tensile test percent elongation.

The above photomicrographs-of FIGURES 1 to 7, in-

clusive, are those respectively taken of a representative portion of the longitudinally cut surface of longitudinally sectional diameter temperature treated test bars of the alloys represented after such cutsurfaces had been mechanically polished and then electrolytically etched for 15 to 25 seconds in a 10%oxalic aqueous acid solution ata current density of between 0.25 to 0.5 ampere per square inch.

The term austenitic alloy as used herein is intended ing elements such as nickeL-nitrogen, manganese, and

carbon andwhich is austenitic at room and service tem peratur'es and will nottr-ansforrn to a bodycentered cubic (B.C.C.) solid solution or .ferritic state upon cooling v room temperature after service at high temperature. I T y As mentionedbriefly above, our.inventionjis, among other things, based upon the discovery that stable austenitic ferrous alloys, especially those having stainless qualities and other combinations of physical properties described above, may be obtained by combining small critical quantities of certain austenite forming elements of the group nickel (Ni), manganese (Mn), carbon (C), and nitrogen (N) of which carbon is essential and is relatively high and nickel is essential except in the case of particular combinations of manganese and sometimes nitrogen, with an iron base material and with small critical quantities of a plurality of-certain NaCl-type carbide formers of the group molybdenum (Mo), tungsten (W), columbium (Cb), tantalum (Ta), titanium (Ti, and vanadium (V), to produce a stable austenitic matrix at room and operating temperatures and to provide a balanced alloy with outstanding combinations of room and elevated temperature properties. Such features make it possible to obtain iron base alloys possessing properties comparable or superior to alloys high in nickel and/or cobalt.

' The new alloys are lowcost, investment-castable, weldable austenitic stainless steels preferably containing about 60% and more iron by weight and having a 100-hour stress-rupture strength at 1500 F. of the order of 22,000-28,000 p.s.i. and even higher in the as cast condition. These alloys also possess high strength at room temperature and are more readily weldable than the known H831 and HS21 alloys.

In general the preferred stable austenitic compositions of our invention of the aforementioned character having one or more of the described properties, and which may be free of or contain small amounts of cobalt will preferably contain in addition to iron as essential ingredients the elements carbon, silicon, nickel, manganese, nitrogen, chromium, molybdenum, tungsten, and columbium and/ or tantalum preferably columbium.

The amounts of each by weight percent will depend upon the specific properties desired and may be broadly expressed by a composition containing as essential ingr'edients at least about 0.8% and up to about 1.25% carbon, about 2 to 8% nickel, about 1 to 15% manganese, about 12 to 35% chromium and a plurality of molybdenum, tungsten and metal from the group columbium and tantalum in amount not greater than about 12%, preferably not more than about 9% with the molybdenum when used within about 0.1 to 9%, tungsten when used within about 0.1 to and columbium and/or tantalum when used within about 0.1 to 5%. In no case should the combined nickel and manganese exceed about The composition will also desirably include silicon in amount between 0 to about 2.5% preferably at least about 0.3% and nitrogen in amount between 0 to about 0.6%, the latter being especially significant when the nickel is at the 2% level or less and especially when the carbide forming elements are above a total of 3% in the latter eases nitrogen additions of 0.05 to 0.6% and preferably between about 0.05 to 0.3% are a decided advantage. The alloy. may also optionally include cobalt in amount between 0 to about 8% in which case the combined total of nickel, manganese and cobalt should not exceed about 15 The remainder of the composition will essentially be iron which should constitute at least 40% and may be as high as 75% of the alloy. Compositions within the foregoing range of ingredients will provide alloys having the physical properties of at least about 2% elongation with a tensile strength of at least about 85,000 p.s.i. at room temperature, a hot tensile strength at 1500 F. of at least about 40,000 p.s.i. with less than 35 elongation, and a 100 hour rupture life at 1500 F. of at least about 20,000 p.s.i. believed a desirable minimum for commercial applications.

6 24,000 to 28,000 p.s.i. will be obtained when the carbon content is above 0.9% and for optimum results above a minimum of about 1% and up to about 1.1%. Furthermore, in certain instances when the optimum properties provided by the presence of both nickel and manganese are not required, manganese may be entirely omitted in the above compositions or used in amount less than 1% providing at least about 2% nickel be used. Moreover, nickel may be used in amount less than 2% providing that between about 5 to 10% manganese is included and at least about 10% where the nickel content is zero. The amount of manganese in these cases may be reduced to as low as about 2% if some nitrogen preferably at least about 0.3 is provided.

In place of a portion of the iron content in the above compositions and while maintaining at least a 40% concentration there may also be included in the alloy in certain cases for the special advantages hereinafter noted they may supply, certain other elements, for example up to 5% titanium, up to 5% vanadium, up to 1% boron, up to 0.2% phosphorous and such incidental elements as zirconium, aluminum, and magnesium which may be present as a result of master addition alloys employed during processing, reactions between melt and crucible, deoxidizers, degasifie-rs and purifiers which do not deleteriously effect the properties of the alloy (the total of which in cidental elements will not ordinarily exceed about 1%.). Hence the expression the balance or the remainder whenever used in reference to the composition shall be construed to include any of these additions and contaminating elements. To avoid the inclusion of uncontrolled amounts of contaminating elements which may cause some degradation of physical properties it is preferred that starting ingredients be employed whose compositions are known both as to primary and trace elements.

We have particularly discovered that a combination preferably in substantially stoichiometrically balanced amounts of the NaCl-type of carbide formers, molybdenum, tungsten and at least one from the group consisting of columbiunrand tantalum, preferably columbium which is found to be superior to tantalum inthe properties given to the matrix, in relatively small properly proportioned quantities, in a suitable austenitic matrix containing between above about 0.9 and up to 1.1% carbon and other essential ingredients 'within the limits previously stated will produce low cost alloys possessing exceptionally high resistance to stresses at high temperature. Based upon experimentation, we preferably maintain the amount of molybdenum within the range of about 0.2% to 7.0% tungsten Within about 0.2% to 6% and columbium and/ or tantalum alone or combined in either case within the range of about 0.2% to 4.0% and with a total of molybdenum, tungsten, columbium and/ or tantalum not exceeding about 9% by Weight.

Thus for making lowest cost precision castable alloys having a tensile strength of at least 85,000 p.s.i. at room temperature and at least about 2% elongation, a hot strength of at least 40,000 p.s.i. at 1500 F. with less than elongation and a 100 hour rupture strength at 1500 F. of at least 24,000 p.s.i., believed adequate for uses which have heretofore required high nickel or cobalt alloys, the alloy compositions will contain as essential ingredients in critical amounts about 0.9% to 1.1% carbon, preferably above a minimum in the'order of 1.0%, 0.1% to 2% silicon, 12% to 25% chromium, 0.2% to 7% molybdenum, 0.2% to 6% tungsten, 0.2% to 4% columbium and/or tantalum, with a total of molybdenum, tungsten, columbium, and/ or tantalum not exceeding about 9%, 2.0% to 8% nickel, 0% to 9% manganese with a total combined maximum content of nickel and manganese in the order of 11%, the balance being essentially iron in amount at least about 40%, preferably and as high as about In place of part of the iron and while maintaining at least a 60% iron concentration, there may be also included 0% to 7 change occurring in microstructure.

a, 1 assoc,

3% titanium, to 3% vanadium, 0% to 0.4% boron,

0% to 0.2% phosphorous, 0 to 0.6% nitrogen and the incidentals mentionedabove'. Cobaltmay be optionally added in amount between 0% to 8% in which event the combined total of nickel, manganese and cobalt should not exceed 15%. i when: the molybdenun1,-tungsten, and columbium (or tantalum) are used in substantially equal amounts in th range of about 1% to 2% of each. Equal amounts are likewise preferred above the 2% level. Where themolybdenum', tungsten, colurnbium (or substituent tantalum) are used in substantially equal amounts at levels from about 2% to 4%, nitrogen in amount between 0.05 to 06% is found beneficial to the mechanical properties of the alloy and should preferably be included,

Carbon must always be present in the new alloys, it not only being an austenite former but more particularly combining with the NaCl-type of cubic carbide formers molybdenum, tungsten, columbiurn, tantalum, titanium, and vanadium with which it must bebalanced, preferably stoichiometrically to form among other carbides thesi-mple carbides MoC, WC, CbC, TaC, TiC, VC and complex carbides of these with chromium such as the M C types. For optimum creep-rupture strength, the iron base alloys of our invention require a carbon content above an amount in the order of about 1% and in the range 09 to 1.1%, preferably stoichiometrically balanced with the addition of these strengthening elements.

Moreover, best results are obtainable forming elements present in the, alloy. 'Multiple addi-' tions of carbide forming elements in optimum amounts impart much greater creep; resistance thanany single addition thereof incomparable or even greater-amount.

Thusthe most'potent strengthening elements are found to be'tungsten, molybdenumand ,co'lumbium. ,We have discovered thatoptimum high temperature properties are i in example No. 6 inthe tables, containing approximately 1% each by Weight of W,'Mo, and Cb, about 1% C,

about 20% Cr, about 5% Ni, about 4.5% Mn, and

about 1% Si displayed a 100 hour rupture strength of 27,000 p.s.i. at 1500 F., a cold tensileductility of 4% elongation, good oxidation resistance upto 1800 F. and

Variations in carbon content not only profoundly aifect v the stress rupture strength at 1500 F, the room temperature percent elongation andthe stress rupture elongation V at 1500 Fjof alloys having compositions otherwise mvithin the foregoing ranges of ingredients, but are found to V ductility in this range, and that the stress rupture elon-' gation at 1500 F. similarly undergoes a surprising change at 0.8% carbon reversing-the expectedcontinued downtrend with increasing carbon content and is in a safe preferred range of between '5 and elongation. More over, it is found that alloys with a carbon content below 0.8% carbon have an appreciable change in microstructure from those in the range of 0.8% to 1.4% carbon and that below about 0.5% carbon there is a drastic These distinctions in structure and physical properties will .be readily evident from the photomicrographs of FIGURES 1 tol 7,.in

, clusive, and the above descriptions thereof, and frornthe graphs of physical properties in FIGURE 8 of a typical I base composition with varying carbon content. From this data it should be apparent that alloys below 0.'8%

carbon have substantially weaker physical properties and that increases in carbon above about 1.25% does not appear to strengthen the alloys to any appreciable extent at least to be of practical value, but renders the cold duca tility ofgt-he alloys in most cases inadequate. Hence carbon in the range 0.8% to 1.25% is preferred and for optimum results carbon above an amount in the order of about 1% and in the range 0.9% to. 1.1% gives best results. a

Thus we have discovered that A further feature of our'invention is the critical proportioning of the NaCl type carbide formers to the carbon content of the alloy. The resulting complex oarbide phases are highly resistant to spheroidization atele vated temperatures and give to the alloy matrix; great .tion is that the alloy be austenitic.

was readily weldable. The tables show examples using other equal amounts of these strengthening elements. See for example, alloys 303, 14, 15, 4, and 7 from which it will be evident that the hourfstress rupture strength is optimum in the range of 0.1% to 2% increasing with quantities up to 1% and decreasing'withincreasing quantities above 2%.

as-an upper limit. v I

Simultaneous variations in the amount of W, Mo,and

Cb at the 6% to 10% levels depressed the stress rupture properties of the basic alloy compositions. However, large additions of molybdenum using smalleramounts of tungsten and columbium within vthe.l2%-total irnprovedthe hot tensile strength while high oolurnbium additions hadthe' reverse effect. Large additions of tung-.

sten acted'similarly to the molybdenum.

.Titanium and vanadium are also possible additives in lieu of or as a replacement of part of the columbium and tantalum, used and preferablyfwithin the limits. set forth for these latter elements. Titanium and vanadium are not especially recommended since experiments seem to indicate that they are proneto permit moresevere oxidation of the matrix during service than where columbium or tantalum are used without them andthus tend to, lower the ultimate stress value. Titanium and vanadium are also not as suitable as columbium and tantalum from, the standpoint of room temperature ductility and alloy castability. They tend to Produce lower ductility and in certain cases aifect mold reaction during investment-type casting,resulting in castings with surfaces not as smooth as those obtained when c olumbium .or tantalum are employed from the stated group. 1

Aspreviouslynoted, an essential feature of our inven- For this purpose the nickel content in the absence of manganese, or nitrogen or cobalt should notbe less than' about' 2%, It is also desired and important that the austenite formers others than carbon, i.e., nickel, manganese, nitrogen, and cobalt should be adjusted to obtain in conjunction with the essential'carbon content an austenitic, composition that is not only stable but of optimum high temperature properties, a condition attained when the non-strengthcn ing austinite formers (nickel and cobalt) are not present in excessive amount, and the austenite formerscarbon, manganese, nitrogen-Which form strengthen'ing carbides and nitrides areoptimum in amount. Hence the amount ofnickel required will. decrease with increasing amount of austenite .formers carbon, manganese, and nitrogen used. The amount of nickel needed'to contribute to.

stability will also vary withthe character and amount of the other essential ingredients oft'he alloys especially A maximum of 4% of each is preferred.

sten, columbium, tantalum, and silicon. For the amounts of carbon used in our alloys, an amount of nickel between about 2% to 8% will produce optimum properties. These properties are far superior to those attainable with high concentrations of nickel and lower concentrations of carbon. In fact, these properties approach, and even surpass those of the known super alloys as will hereinafter be evident.

Where cost is not material and optimum high temperature properties are not essential, higher nickel concentrations than 8% may be used in conjunction with the other essential ingredients of the alloy to obtain some improvement in oxidation resistance. Amounts of nickel in excess of 20%, for instance up to 35%, will not only adversely affect room temperature and high temperature stress rupture properties but will also decrease the solubility of nitrogen in the alloy to a point where additions of the latter in amount to provide desirable properties thereby in the alloy will be inhibited. Moreover, such may require a reduction in the minimum amount of iron desired in the alloy and for many applications will make cost prohibitive. Hence, nickel in amounts up .to about 8% is preferred and amounts in excess of 20% not generally advocated.

Manganese is also preferably employed in producing the alloys of our invention and in amounts up to about 15 by weight. Amounts between to 7% are preferred as in this range, the manganese in addition to being a nitride former, is also a relatively strong austenitizer and is somewhat effective in improving the casta'bility of the alloy. About gives best results. Replacing nickel with manganese or raising the manganese at the standard nickel level of about 2% to 8% will produce higher tensile strentghs at 1500 F. with no apparent change in rupture strength. The cold tensile ductility may in certain cases, however, be somewhat below optimum.

The addition of nitrogen to. the alloy composition is sometimes advantageous. It may be used as a substitute for other austenite formers and when included tends to stabilize the austenite composition. Moreover, nitrogen combines to some extent with the chromium of the alloy and to a smaller extent with the manganese when the latter is present to form nitrides and with columbium to form carbonitrides. In certain alloys the production of precipitation hardening nitrides is believed to be partly responsible for the improved high temperature strength and the high temperature service life of the alloys of our invention. Many of the better alloys produced by our invention will contain optimum amounts of nitrogen. The total amount of nitrogen that can be held in stable combination in the alloys depends upon the nickel, chromium, columbium and manganese content of the alloy. Nickel as previously noted when used in large amounts, reduces the amount of nitrogen that may be present. Chromium and manganese have the opposite effect. When nitrogen and manganese are used in optimum amounts as in alloys 86a-86d for example, it is even possible to obtain an alloy with excellent properties without using any nickel or cobalt. Too high uncombined nitrogen is believed to be detrimental to the high temperature properties of the alloys since higher porosity of the casting is often observed where too much nitrogen is in the mass. This element will preferably be used in amounts not exceeding about 0.6%.

As previously described, the alloys of our invention are preferably endowed with stainless properties. For this purpose chromium is an essential ingredient; a minimum of about 8% and a maximum of 35% is preferred. We have found that amounts between about 12% to 25% usually provide adequate oxidation and corrosion resistance at high temperatures up to about 1800 F.

Silicon usually occurs as an impurity in the alloys of our invention in amounts up to about 0.5% to 0.7% (see tables). This is especially the case where commercial grades of master alloys are used in their preparation. The

v strength.

. l0- effect of silicon is threefold. First it influences the solidification characteristics of the alloy such as fluidity, caste ability, and weldability. Secondly, it influences the phys ical properties at both room and elevated temperatures. Thirdly it is a strong deoxidizer.

Thus with regard to weldability, the presence of at least 0.4% silicon in the stable austenitic alloys of our invention, lean in austenite formers, provides a balanced structure with the embrittling carbide forming elements when the carbon content is above 0.8% whereby there is present in the network a relatively small amount of dispersed ferrite. This facilitates the making of a ductile weld.

It is known that castable iron alloys with over 0.8% carbon are not readily weldable. Also that most super alloys are so brittle with even much less carbon that they must not only be preheated to prevent cracking during welding but must also be postheated to relieve stresses produced during welding, to avoid after cracking. The alloys of our invention avoid these steps and conditions so long I as the amount of silicon present remains under about 2.0%. Silicon above this amount will embrittle the alloy particularly at room temperature and reduce the stress rupture strength at 1500 F. Moreover, silicon is a very effective deoxidizer and as such not only influences the ductility of the alloy as prepared, but also the character of any weld. Furthermore, it improves resistance of the alloy to oxidation at elevated temperatures.

The iron base alloys of our invention may also include certain other specific elements such as phosphorous, boron and sulphur. Thus as previously indicated, boron may be added in amounts from 0 to 1% to improve the high temperature strength of the alloy. It is found, however,

that boron in amounts above 0.04% will cause a decrease in room temperature ductility. Phosphorous when added in amounts up to about 0.1% renders the alloy susceptible to heat treatment with corresponding improved hot Increasing amounts of phosphorous tends to decrease the ductility of the resulting alloy.

Although the alloys of our invention are characterized as being cobalt-free, additions of cobalt are found in some cases to have a beneficial influence on the castability and the high temperature properties of the alloys and where some additional cost is warranted between 0% to 15% of this element may be included in the alloy composition.

Iowever, it is not required.

A composition range Within the above broad limits that has been found to be particularly adaptable for making commercial alloys comprises alloys containing chromium in amount between 18% to 22% by weight, carbon 0.8% to 1.25%, silicon in amount between 0.1% to 2%, molybdenum in amount between 0.5% to 3%, tungsten in amount between 0.5 to 3%, columbium in amount between 0.5 to 2%, nickel in amount between 2% to 6%, manganese in amount up to 7%, nitrogen in amount up to 0.3% with the balance essentially iron (Fe) and any incidental impurities. Alloys within these commercial ranges are machinable and readily welded and have remarkably high strengths and stability at high temperature up to 1800 F. v

In evaluating heat resisting alloys for such application as gas turbine engines and the like, long time rupture tests at high temperatures and short-time tensile tests at room temperature and at elevated temperatures are desirable indicators of their efiectiveness. Such tests have been made of the novel alloys of this invention using precision cast test bars, produced by the well known lost wax process, measuring about 0.25 inch in diameter and having a one-inch gauge length. For determining the suitability of the alloy for high temperature application, a long time test, so-called stress rupture test was used. In this test each of several specimens of the alloy was sub-' jected to measured tensile stress at a particular elevated temperature, and the time required for the sample tofail under these conditions of temperature and stress was noted. Data obtained in this manner permits determination of a See footnotes at end of table.

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LLLLLLLLQJLLLLLLLQQLLLLQLLLLLLLL WQQQLQQQQQQQQLLQLQQLLLLLLLLLLQQLQQLQLLLLLLLQQLLLLLLLQL Alloy ally equal parts bf rare earth metals such as Ge, La)

l6 Made from elemental hi 15 10; 1 TABLE NO. II 7 TABLE NO. IIContznued T st e Dam Short-Time Tensile Testing Stress-Rupture, Testing at Short-Time TensileTesting Stress-Rupture Testing at 5 Room Tem- 1510 F;

Alloy peratme r 1 Stress. Time, Per-. 100,111. o 1.501? (p.S.i.) Hour cent Life Alloy perafiure TS 1901- TS Per- E1 (p.s.i)

S9955 T1199. 1001119 .31 cent (p.s.i.) m:

(p.s.1.) Hour cent L1f e EL E1; TS Per- 'IS Per- El. (p.s.1.) 10 (p.s.1.) cent (p.31) cent 95,000 5 40,000 22 22,500 7 5 20 30,000 17 3.-- 98, 400 3 47,000 20 20,000 477 7.0 90 300. 2 50,600 17 22,500 144 000 8 5 49, 100 17 25, 000 41 3A 107,000 4 50,200 15 20,000 604 10.5 31 93,000 1 ,500 20 20,000 110 25, 000 04 13. 0 24, 000 57, 000 24 22,500 21 3B 108.000 2.5 50,750 15 20, 000 470 4.0 32;--- 32,400 1 43,200 15 22,500 245 000 77 13. 5 000 102, 300 2. 5 43, 300 15 25, 000 123 3C 90,000 1.5 52,500 13 20, 000 455 21 45,400 13 27,500 20 ,000 153 9 30- 30,000 1 72,250 22,500 27, 500 45 14 02, 250 1' 25,000 42 32,500 43 5 1 r 30,000 19 31,000 54 4 37 119,250 1 01,500 15 20, 000 413 30,000 105 4.0 119,700 1 00,500 13 22,500 130 0. 105,000 4 41,000 13 20,000 1,215 7 53,750 19 25,000 71 5 25, 000 0 7. 5 1 l 27, 500 29 5 27,500 125 0 7 33.--- 114,100 r 1 01,000 13 20,000 033 100,000 1 45,000 3 20, 000 15 45 110,000 1 00,300 17 22,500 130 3. 95,000 2 44,000 13 20,000. 193 0 00,200 13: 25,000 71 5 See footnotes at end 0! 1:31:10

TABLE NO. II-Cntznued TABLE NO. IIC0ntmued Short-Time Tensile Testing Stress-Rupture Testing at Short-Time Tensile Testing Stress-Rupture Testing 2t Room Tem- 1,500 F. Room Tem- 1,500 F. Alloy perature Alloy perature Stress Time, Per- 100 Hr. Stress Tim e, Per- 100 Hr, (p.s.1.) Hour cent Life (p.s.i.) Hour cent Life TS Per- TS Por- E1. (p.s.i.) 10 TS Per- TS Per- E1 (p.s.i.) (p.s.i.) cent (p.s.1.) cent (p.s.i.) cent (p.s.i.) cent E1. E1. E1. E1.

46, 800 5 25, 000 65 6. 5 10,000 6 44, 550 10 27,500 10 98 123,000 3 53, 000 15 20,000 238.5 5 66...- Brittle 76, 000 2 25, 000 334 14. 5 25,000 29 7 74,100 3 27,500 146, 5 18 99...- 113, 300 3 42, 700 14 20, 000 0. 2 42 67..-- 126, 600 1 66, 500 9 22, 500 689 6. 5 10,000 6. 5 48 63, 000 13 25, 600 258 12 7, 500 55. 4 26 27, 500 140. 5 15 100- 85, 090 1 21, 000 20, 000 0.2 68..-- 123, 600 2. 5 64, 200 13 22, 500 389 2. 5 10, 000 7. 0 42 65, 000 16 25, 000 127 15 7, 500 36.8 24 27, 500 58 10 2 41, 750 14 20,000 19 11 30, 000 27 15 17,500 44 7 69..-- 101, 000 1 73, 000 3 25, 000 148. 5 22. 5 15, 000 162. 5 8

See footnotes at end of table.

1 9 TABLE NO. II-Com'inued Short-Time Tensile Testing Stress-Rupture Testing at Room Tem- 1,500 F. Alloy perature Stress Time, Per- 100 Hr. (p.s.i.) Hour cent Life TS Per- TS Per E1. (p.s.i.)

(p.s.i.) cent (p.s.i.) cent El. E1.

1 Solution heat treated condition, 2,150 F.l1 hr.+oil quench.

1 Stress rupture tested at 1,700 F. 3 Tested at 1,350J F.

Tested at 1,200 F.

5 Tested at 1,650" F.

6 Tested at 1,800 E. v

Ineach example given in the tables the alloy was'prepared either from relatively pure starting ingredients or from commercial grade; source materials such as electrolytic grade iron for'iron, 'ferro alloys of chromium, molybdenum, tungsten, and .columbiurn for each of these metals, commercial grades of electrolytic nickel and elec-. trolytic manganese for these substances and high carbon ferro-ch rome as a source of carbon. In some instances the ironfcomponent was derived from SAE 1010 steel or Arm'co iron. Melting of proper proportions of the ingredients was normally carried out in a magnesia crucible and sometimes in a zircon or zirconia crucible and in each case in a high frequency induction furnace with the melt protected by an argon blanket or under a blanket of slag. I

As previously notcd that cast alloys may contain. contaminating elements present by reason of using commercial alloy starting materials, by pickup from the crucible, etc. possess physical properties withinthe limits given for a Although all of the alloys produced will alloy composition of the essentialingredients these contaminants may to some degree aifect the hour stress rupture values, such depending onthe batch of commercial grade materialsemployed and the kind and number of contaminants contained therein. Such effects may understood by a comparison for example of alloy Nos.

6 and 302, the former beingmade using relatively pure for in this way the compositions may becontrolled and deviations in properties explained.

Melts of the alloys recorded in the tables were cast in standard investment molds or ceramic shell molds produced by the lost wax process to obtain precision cast test bars. Each cast alloy was subjected to the stated short time tensile testing at room temperature and at 1500 F. and at a plurality of stress levels. The results of these tests recorded in Table I1 demonstrate the outstanding characteristics of the alloys of our invention.

The effect of varying the concentration of the essential ingredients of the alloys of our inventionare evident from the many examples given in the tables. For the purpose of discussing some aspects thereof reference'will be made to the efiects obtained on a composition containing 20% Cr, 5% Ni, and 5% Mn. It will be understood that such is not to be construed as limiting such alloys to these specific values for other heats show the efiects onv other compositions.

Thus the effect of varying carbon additions on the properties of such a basic alloy is to be observed for example from the composition and test data for the following groups of heat numbers and from the graphs in FIG- URE 8. Thus heats 124, 123, 120, 6, 104, and 119 show 'the effect of varying the carbon content using a 1% level of W,'Mo and Cb and approximately 0.5 Si. Heats 1, .2, and 15 show the efiect at levels of 2% each of W, and Mo and 1% Co and 0.5% to 1% Si. Heats 91, 13, and show theefiect of carbon at similar levels of W, Mo, Cb and Si where about 0.1 N is'included. Heats 77, 76, 15, and 117 show the effect of carbonat levels of 2% each of W, Mo, and Cb, 0.5% Siand 0.3% N. Heats 45, 62, 66, and 69 show the effects of varying carbon at a 4% level of'W and Mo, 1%of Cb, 0.5% Si, and 0.3 N.

This data reveals that the optimum high temperature properties are achieved'at critical levelsof carbon. For instance, lowering the carbon content substantially below about the 1% level for instance below 0.8% causes considerable .weakening of the alloy whereas in general raising such amount substantially above the 1.25% level does not appreciably strengthen the same. It. should also be noted that the optimum level of carbon increases with increasing concentrations of the NaCl-type of carbide forming elements. Moreover, those alloys containing amounts of carbon substantially above 1.25 exhibit low and inadequate cold ductility for many applications. Hence any potential improvement in strength due to higher carbon concentrationssubstantially above 1.25 appears to be without substantial practical value.

Asdescribed above, certain advantages result from the presence of nitrogen in the austenitic alloys of-ourinvention.. The effect of various concentrations up to 0.7% thereof is apparent, for example, from a consideration of the mechanical properties of an alloy composition containing 20% Cr, 5% Ni, 5% Mn, 1% C, and 0.5% silicon at diflierent levels of W, M0, and Cb as recorded for heat Nos. 3, 12, 13, 89, 14, 15, 106, 64 when 2% each of W, Mo and Ch were used. Heats 18, 26,67, and 116 show thecifect at the same level of the carbide formers but at 7%1evel of Mn and 4% Ni.' Heats 6,102,90, and 92 show the effect at a 1% level of W, Mo, and Ch. Nitrogen is an effective austeniteformer and a potent strengthener through its formation of simple and complex nitrides. Significantly it improves the hot strength of the composiates o tion in the range 0.05% to 0.3% and without apparent loss in cold tensile ductility. Alloys 12-15 displayed outstanding creep strengths at 1500", these registering stress rupture values of 27,000 psi. after 100 hours. The best stress rupture value at 100 hours however with some sacrifice in the cold ductility properties was given by alloy 67 in which the nitrogen concentration was 0.5 This alloy contained as noted, 4% Ni and 7 Mn. It had a strength of 28,000 psi. Alloy 12 on the other hand containing 0.05% nitrogen and 5% each of Ni and Mn had a 100 hour rupture'strength of 27,500 psi. with good cold ductility. Somewhat lower ultimate creep rupture strengths were obtained using amounts of nitrogen in the same range when the W, Mo and Ch concentration was lowered from 2% to 1% as in alloys 6, 102, 90, and 92. In these instances however, the hot tensile strength was increased with increasing quantities of nitrogen. Alloy 195 containing 0.3% nitrogen, 5% manganese, and 2% of each of W, Mo, and Cb showed exceptionally good properties in the absence of both nickel and cobalt.

The effect of varying amounts of chromium in a basic composition containing for instance 5% 1i, 5% Mn, 1% C, and 0.5% Si were considered. The tables indicate that best results are to be expected from concentrations between about 12% to of chromium at these levels of austenite formers. At chromium the creep rupture strength dropped materially. Apparently this alloy becomes structurally unstable at this concentration. See and compare, for example, alloys 16, 10, and 17. Stable alloys using larger amounts of chromium than 20% can be obtained as evident from a consideration of heats 40, 39, 38, and 41 by employing greater amounts of austenite formers, in this instance nitrogen being used.

The ei fects of varying the total concentration of the carbide formers W, Mo, and Cb when used in multiple in the ratio of 1:1:1 will be seen from a comparison of the properties of alloys 129, 94, 6, 3, 4, 7, and 303 in connection with a composition containing 20 Cr, 5% Ni, 5% Mn, 1% C, and 0.5% Si.

In this illustration best properties were obtained at the 1% level of the carbide formers. When nitrogen in amounts of about 0.1% were added as in heats 93, 90, and 13 the best results were obtained at a higher level of the carbide formers W, Mo and Cb for instance 2% each. It will be observed that the best combination of properties is produced by a composition such as alloy No. 6 corn still using them in multiple had different results. Alloys" 6', 29, 30, 31 using a 1% level of each of W and Cb indicate some slight lowering in general of stress-rupture strengths at 100 hours for increasing amounts of Mo. Alloys 3, 21, and 22 show the effect of increasing amounts of M0 at the 2% level of each of W and Cb. Alloys 57, 56, 46, and 45 show the effect where unequal and greater amounts of W and Cb were used for instance 4% W, 2% Cb and 0.3% N was present. The same was true of tungsten with fixedequal and unequal amounts of Mo and Cb as demonstrated by alloys 15, 46, and 47 employing 2% each of Mo and Ch andalloys; 58, 59, 54, 60, and 45 containing 4% Mo and 1% Cb. Similar results were obtained by varying the columbium with fixed unequal concentrations of Mo and W as seen from alloys 29, 32, and 35.

In the case of the specific examples set forth where the individual concentrations of Mo, W and Cb were quite high, 6% and more, there were marked drops in stress rupture strength apparent but increases in the hot tensile strength properties.

The addition of fourth and fifth NaCl carbide forming strengthening elements, for example, titanium (Ti) and vanadium (V) on compositions of the invention is somewhat evident from alloys 43, 50, 52, 51, and 53. Thus increasing concentrations of titanium progressively lowered the hot strength properties andincreased amounts of vanadium lowered the long time stress-rupture values and the hot strength at 1500 F. Such additions moreover produced alloys more diificult to cast. Moreover, mold reactions in preparation and excessive oxidation at 1500 F. were observed, such apparentlyaccounting in some degree for their lower high temperature strengths.

Use of less than about 2% nickel in the absence of adequate manganese produces alloys having low stress-rupture strengths, due to structural instability of the austenite matrix. On the other hand, although substantial amounts or" nickel up to about 20% may be used, additions of nickel substantially above the 2% to 8% level do not produce si nificant benefits. These results may be noted, for example, from the properties of alloys 86, 118, 85, 84,

. 83, 82, 81, and 78 at the level of 0.3% ,N with Mn omitted. Alloys 75, 79, and 304 show the effects where amounts of Ni from 0% to 7% were used, alloys 75, 79 and 80 comprising alloys containing 10% Mn, 0.5 N, and 0.1% P and alloy 304 containing no manganese or nitrogen.

Theeir'ect of silicon (Si) upon the mechanical properties of a composition having 20% Cr, 5% Ni, 5% Mn, 2% Mo, 2% Cb, and 1% C is evident from alloys 3, 23, 24, 25, and 125. These results show that up to about 2% silicon does not adversely affect mechanical properties and a 3% level apparently causes embrittlement at room temperature and a low stress rupture strength at 1500" F. Silicon as described above, is desirable in some small amounts as a deoxidizing agent and for improving castability and weldability of the alloy.

Phosphorous is not required. However, when used in amount up to 0.1% and in the absence of nitrogen, improves the room temperature and high temperature properties of the alloy in the as cast condition. This will be evident by a comparison of the properties of alloys No.

3, 9, 10, and 8 and of alloys 5 and 13.

Additions of boron in amounts up to 0.2% as shown by alloys 3A, 3B, 3C and 31) appear to improve the stressrupture strength at 1500 F. as evident from a comparison with alloy 3. Above this concentration with this base composition no further improvement was evident. Moreover, increasing amounts within the range up to 0.2% gave a slight improvement in hot strength but recorded a substantial drop in cold ductility when above 0.02 concentration, the elongation then decreasing from 4% to 1% when the boron content was 0.2. Addition of 1.0% boron in compositions having a lower total of carbide formers as in example 136 also improved the stress rupture strength but produced a lower value of cold ductiiity.

Out of the many alloys made in accordance with the teachings of our invention and for which data is recorded in the above table, those preferred by reason of high temperature strength, and room temperature properties are found, for example, in alloys 6 and 15. Alloy 6 in addition to iron and impurities contained substantially 20% Cr, 5% Ni, 5% Mn, 1% Mo, 1% W, 1% Cb, 1% C,-and 0.5 Si, and alloy 15 similar amounts of Cr, Ni, Mn, C, and Si, about twice as much of each of Mo, W, and Cb and including between 0.05% to 0.3% N.

These two alloys have been subjected to extensive tests to determine the following properties:

(1) Short time tensile Strength at temperatures up to 1800 F. and stress rupture and creep data at temperatures of 1200 F. to 1500" F.

(2) Thermal shock behavior as compared to cobalt based so-called super alloys.

(3) Thermal expansion characteristics.

(4) Oxidation behavior.

(5) Microstructural characteristics in the ascast condition.

(6) Welding characteristics under restrained and unrestrained conditions.

The short time tensile data is given below in Table lll:

TABLE 111 Short Time Tensile Properties Alloys 6 and J Alloy 6 Alloy Testing Temp.

UJIXS. Percent Percent U.T.S. Percent Percent (p.s.i.) E R. (p.s.i.) E1 RA.

Room Temp--- 105,000 4. 0 2.0 125,000 3 0 2.0 600 F 86,000 5.0 2.0 78,000 5.0 2. 5 93, 500 5.0 4. 6

From this short time tensile data it is apparent that v alloy 15 is much stronger than alloy-6 over the temperature range recorded. On the other hand, alloy 6 is possessed of a higher degree of ductility at room tempera ture.

The stress-rupture strengths for alloys 6, l2, and 15 and of the cobalt base alloys H821 and H331 at elevated temperatures are given in Table IV below:

TABLE IV Stress Rupture Characteristics of Alloyso and 15 at High Temperature as Compared to Cobalt-Base Alloys H531 and HSZI 1 Values for HSBI and H821 as reported in Haynes A1l oys, published by Haynes Stellite Company (Division of Union Carbide and Carbon Corporation).

Z Heat-treated condition.

It'will be apparent. that the creep-rupture strengths of iron base alloys Gand 15 are much superior to the cobalt base alloy HS2lover the entire range of temperatures tested, and superior to H831 at temperaturesup to l400 F; and substantially equivalent in the temperature range of 1400 F. to 1500 F.

The oxidation resistance of alloys 6 and 15 and cobalt alloy H331 were evaluated at temperatures of 1500 F,

. E l 1 I 1800 F., and 2000 F. Cyclingoxidation tests as'well as continuous exposures in still air were made. The re-. sults based upon weight lossindicate' that both of alloys 6 and 15 have comparable refractory characteristics'at elevated temperatures up to about 1500 F: but their oxidationresistance is not as great as cobalt base alloy H831 at 1800 F. and above. The scaling temperatures of these iron base alloys 6 and 15 is in theorderof 1800 F. which is comparable to other types of stainless steels containing about 20% chromium. V p

The welding characteristics of iron base austenitic alloys such as alloys 6 and 15 indicate many advantages. Studies show that sound and good weldments can be obtained in both restrained and unrestrained conditions and without the preheating and post heating: treatments required on other alloys such as H831. The latter cobalt base alloys must be preheated to about 600 F. prior to Welding. Unless this is done, stresses created provide cracking in the joints and'other portions of the welded parts. After welding H331, parts must be post heated to 1200" F. to release welding stresses.

Alloys of our invention such as alloys 6 andlS do not require these operations to be welded. No adverse ei'lects are noted by visual, X-ray and radiographic examinations where such steps were eliminated. Al though the reason therefor is not definitely known, it is believed that the pesence of silicon in these alloys accounts to a substantial degree for their improved weldability and improved ductility.

Alloys 6 and 15 have been welded in the unrestrained condition using commercially available alloy welding rods such as l9/9WMo (a stainless steel base alloy), l -665. (a cobalt basealloy) and .N-lSS (a cobaltnickel rich iron base alloy). -Turbine wheel ring blades of alloy 6 were welded (a restrained weld) to a 16-25-6 stainless steel alloy hub using a 19/9 WMo filler rod and by operations using a 600 preheat and 1200 F. post treatment and by operations omitting these preheat and post heating steps with equal efiect.

The welds are good. There is no tendency toward cracking of these alloys 6 and 15 duringwelding in a estrained condition.- This is especially significant.

Moreover, the material ilowswell and showsfa .good heat-affected zoneimicrostructurally. There was little if any difference noted in welding with various frods used for unrestrained welds; i i

The iron base alloys of our invention, as previously described, are austenitio. They furthermore contain in addition to other types of compounds which are also precipitated in the solid solution matrix, one orjmore of the primary or massive, ledeburite or eutectic and spheroidal or dot-type types of carbides.

preferred alloys, show a much more balanced alloy struc ture than those discussed above in terms of.distribution and relative amounts of secondary phases, Our studies also make it appear that optimum stress-rupture strength in the alloys of our invention is dependent upon a micro? structural equilibrium betweensolid solution matrix composition, primary and eutetic carbides, available carbon and carbide forming elements and available nitrogen and nitride formers for matrix strengthening as well as proper Photornicro- V graphs of metallographic specimens 'of: a compositionsuch as alloy 3 showedthat this alloy ingthe ,as-cast conditionconsists of an austenitic matrix embedding dot- 

1. AN IRON BASE CASTING TYPE ALLOY FOR PRODUCING PRODUCTS OF COMPLEX SHAPE AND HAVING SUBSTANTIAL OXIDATION RESISTANCE AND HIGH STRENGTH AT ELEVATED TEMPERATURES, SAID ALLOY CONSISTING ESSENTIALLY OF BY WEIGHT PERCENT, ABOUT 0.8 TO 1.6% CARBON; ABOUT 12 TO 35% CHROMIUM; BETWEEN 0 TO 2.5% SILICON; BETWEEN 0.2 TO 12% OF A PLURALITY OF CARBIDE FORMING ELEMENTS FROM THE GROUP CONSISTING OF TUNGSTEN IN THE RANGE 0.1 TO 10%, MOLYBDENUM IN THE RANGE 0.1 TO 9%, COLUMBIUM AND TANTALUM, SAID COLUMBIUM AND TANTALUM COMBINED BEING IN THE RANGE 0.1 TO 5%; UP TO ABOUT 15% OF METAL FROM THE GROUP CONSISTING OF, NICKEL IN THE RANGE 0 TO ABOUT 8%, MANGANESE IN THE RANGE 0 TO 15%, AND COBALT IN THE RANGE 0 TO 8%; AND FROM 0 TO 0.6% NITROGEN; THE BALANCE ESSENTIALLY IRON, SAID IRON CONTENT BEING AT LEAST 40%, SAID NICKEL BEING AT LEAST ABOUT 2% WHEN SAID MANGANESE IS LESS THAN ABOUT 1%, SAID MANGANESE BEING AT LEAST BETWEEN 2 TO 10% WHEN THE NICKEL IS LESS THAN 2% AND NITROGEN IS PRESENT, AT LEAST BETWEEN ABOUT 5 TO 10% WHEN SUBSTANTIALLY NO NITROGEN IS PRESENT AND THE NICKEL IS LESS THAN ABOUT 2%, AND AT LEAST ABOUT 10% WHEN THE NICKEL IS ZERO; SAID ALLOY BEING AUSTENITIC AT ROOM TEMPERATURE AND BEING CHARACTERIZED IN THE AS CAST CONDITION BY A 100 HOUR RUPTURE STRENGTH AT 1500*F. OF AT LEAST ABOUT 20,000 P.S.I. AND SAID ALLOY BEING CASTABLE INTO COMPLEX SHAPES AND PRODUCTS, A TEST BAR THEREOF 1/4" IN DIAMETER EXPOSED TO 1500*F. FOR 100 HOURS BEING GENERALLY CHARACTERIZED BY A STRUCTURE HAVING AN INTERDENDRITIC NETWORK OF SUBSTANTIALLY LEDEBURITE PHASES OUTLINING AN AUSTENITIC PHASE EMBEDDING RELATIVELY FINE DOT-LIKE PRECIPITATES RANDOMLY DISTRIBUTED THEREIN. 